Effects of Annealing on the Martensitic Transformation of Ni-Based Ferromagnetic Shape Memory

20 Pages·2015·2.47 MB·English

Metals2015,5,484-503;doi:10.3390/met5020484
OPENACCESS
metals
ISSN2075-4701
www.mdpi.com/journal/metals
Article
Effects of Annealing on the Martensitic Transformation of
Ni-Based Ferromagnetic Shape Memory Heusler Alloys
and Nanoparticles
TinaFichtner1,ChanghaiWang1,AleksandrA.Levin1,GuidoKreiner1,
CatalinaSalazarMejia1,SimoneFabbrici2,3,FrancaAlbertini2 andClaudiaFelser1,*
1 Max-Planck-InstituteforChemicalPhysicsofSolids(MPICPFS),NoethnitzerStr. 40,
DE-01187Dresden,Germany;E-Mails: tina.fi[email protected](T.F.);
[email protected](C.W.);[email protected](A.A.L.);
[email protected](G.K.);[email protected](C.S.M.)
2 InstituteofMaterialsforElectronicsandMagnetism,NationalResearchCouncil(IMEM-CNR),
ParcoAreadelleScienze37/A,I-43124Parma,Italy;E-Mails: [email protected](S.F.);
[email protected](F.A.)
3 LaboratoryofMicroandSubmicroEnablingTechnologiesforEmilia-Romagna(MISTE-R),
viaGobetti101,I-40129Bologna,Italy
*Authortowhomcorrespondenceshouldbeaddressed;E-Mail: [email protected];
Tel.:[phone];Fax: [phone].
AcademicEditor: KurtR.Ziebeck
Received: 20February2015/Accepted: 12March2015/Published: 25March2015
Abstract: We report on the effects of annealing on the martensitic phase transformation in
the Ni-based Heusler system: Mn Ni Sn and Mn Ni Sn powder and Co Ni Ga
50 40 10 50 41 9 50 21 32
nanoparticles. For the powdered Mn Ni Sn and Mn Ni Sn alloys, structural and
50 40 10 50 41 9
magnetic measurements reveal that post-annealing decreases the martensitic transformation
temperatures and increases the transition hysteresis. This might be associated with a
release of stress in the Mn Ni Sn and Mn Ni Sn alloys during the annealing process.
50 40 10 50 41 9
However, in the case of Co Ni Ga nanoparticles, a reverse phenomenon is observed.
50 21 32
X-ray diffraction analysis results reveal that the as-prepared Co Ni Ga nanoparticles
50 21 32
do not show a martensitic phase at room temperature. Post-annealing followed by ice
quenching,however,isfoundtotriggertheformationofthemartensiticphase. Thepresence
of the martensitic transition is attributed to annealing-induced particle growth and the stress
introducedduringquenching.
Metals2015,5 485
Keywords: phasetransitions;magnetism;X-rayandelectronscattering
1.Introduction
Heusler alloys that belong to the ferromagnetic shape memory alloys (FSMAs) have been a
fascinating class of materials since they were reported by Ullakko et al. [1] two decades ago.
FSMAs are ferromagnetic and show a magnetic field-induced martensitic transformation. Most shape
memory Heusler alloys contain Ni. Those with Mn and Ga, In, Sn and Sb [2–6] are the most
investigated compounds. They are characterized by large magnetic field-induced strain [7,8], large
magneto-resistance[9,10]andasignificantmagnetocaloriceffect[11–13].
Mn-rich Ni- and Mn-based Heusler alloys have recently become the focus of research owing to
their high localized magnetic moment of 4.0 µ at the Mn atom and their potential for enhanced
B
magnetocaloric properties [11,14]. Mn-rich Mn Ni Sn Heusler compounds were reported to show
50 50−x x
a significant exchange bias field [15], large magnetic entropy change [16,17], small thermal hysteresis
andlargemagnetizationchangeatthemartensitictransformation[18]. However,littleattentionhasbeen
paidtotheeffectofannealingonthosecompounds.
The influence of post-annealing on the martensitic transformation of Ni-rich Ni-Mn-Sn alloys has
been studied by Schagel, McCallum and Lograsso [19] on bulk material and Xuan et al. [20] on
ribbons. Schagel,McCallumandLograsso[19]reportedthatthemagnetizationchangeatthemartensitic
transformationisdependentontheannealingtime. Afteranannealingtimeofapproximatelyfourweeks
at 1223 K, they observed the highest magnetization change during the martensitic transformation and
no overlap of the transition temperatures. Xuan et al. [20], on the other hand, reported that annealing
of ribbons resulted in an increase in the reverse martensitic transition temperature (T ) from 230 K
A
(melt-spun) to 265 K (annealed at 1173 K). Concomitantly, the Curie temperature of the austenite TA
C
increasedfromTA =270K(melt-spun)toTA =275K(annealedat1173K).Thereasonfortheincrease
C C
ofT andTA isattributedtostressandstructuralrelaxationduringpost-annealing.
A C
Inadditiontothesewell-investigatedNi-andMn-basedcompounds,otherFSMAshavebeenstudied,
i.e., Ni-Fe-Ga [21–23], Co-Ni-Al [24,25] and Co-Ni-Ga [26–29]. The Co-Ni-Ga Heusler system was
intensively studied as a promising alternative to Ni-Mn-Ga alloys, especially for high-temperature
shape memory device applications [30–37]. Co-Ni-Ga compounds generally exhibit a dual-phase
microstructure: a parent phase (β) and a non-transformable secondary phase (γ). The ductile Ni-rich
γ-phaseenhancestheoverallworkabilityoftheCo-Ni-Gaalloys. Owingtotheabsenceofamodulatedor
layered structure in martensite, the magnitude of the magnetic field-induced strain (MFIS) in Co-Ni-Ga
alloys is normally much smaller than that of Ni-Mn-Ga alloys [38]. Even so, Co-Ni-Ga compounds are
stillpromisingcandidatesasclassicalthermally-triggeredshapememorymaterials.
TheeffectsofannealingonthestructuralandmagneticphasetransformationofCo-Ni-Gacompounds
have been reported for bulk and ribbon samples [27,39,40]. Liu et al. [27] studied the effects of
annealing temperature and cooling conditions on the structure and martensitic transformation of bulk
Co Ni Ga . It was found that the martensitic transformation temperature increased monotonically
46 27 27
Metals2015,5 486
with increasing annealing temperature and quenching rate. The precipitation of the non-transformable
γ-phase can be avoided if the annealing temperature is higher than 1393 K. Sarma et al. [39]
observed a similar correlation between the annealing temperature and the martensitic transformation
in Co Ni Ga ingots. It was also found that a higher annealing temperature significantly increased
47 23 30
the Curie temperature and moderately decreased the amount of the γ-phase. Dai et al. [40] found
that annealing of the Co Ni Ga ribbons at 673 K suppressed the martensitic transformation and
50 20 30
significantly influenced their magnetic properties. Post-annealing at higher temperatures (>873 K) and
quenching recovered the martensitic phase transition of the Co Ni Ga ribbons. This was attributed
50 20 30
to the elimination of the precipitates and weight fractional variations of the transformable β-phase and
thenon-transformableγ-phasebythepost-annealing.
To our knowledge, no study on the effect of annealing on the martensitic phase transformation
of nanostructured Co Ni Ga nanoparticles, and Mn Ni Sn and Mn Ni Sn powder has been
50 21 32 50 40 10 50 41 9
conducted. Herein,wereportonthesynthesisandannealing-dependentmartensiticphasetransformation
ofCo Ni Ga nanoparticles,andpowderedMn Ni Sn andMn Ni Sn alloys.
50 21 32 50 40 10 50 41 9
This paper is divided into two parts. First, we present the results for the Mn Ni Sn and
50 40 10
Mn Ni Sn powder, followed by the results for the Co Ni Ga nanoparticles. In the following, the
50 41 9 50 21 32
as-preparedpowderisthatwhichresultedfromthemeltingoftheelements,performingaheattreatment
with quenching and grinding in a mortar. The post-annealed powder is that subjected to a subsequent
annealing procedure without quenching. CNG_AS denotes the as-prepared Co Ni Ga nanoparticles
50 21 32
andCNG_AQtheannealedandquenchedCo Ni Ga nanoparticles.
50 21 32
2.ResultsandDiscussion
2.1. Mn Ni Sn andMn Ni Sn Powder
50 40 10 50 41 9
2.1.1. MicroscopicAnalysis
In order to investigate the effect of annealing on the martensitic transformation in Mn-rich Ni-based
FSMAs, the alloys Mn Ni Sn and Mn Ni Sn were synthesized and characterized by optical and
50 40 10 50 41 9
scanning electron microscopy (SEM) with energy dispersive X-ray spectroscopy (EDXS). Figure 1
shows the bright field and back-scattered electron micrographs of Mn Ni Sn . Both methods exhibit
50 40 10
a matrix with an embedded secondary phase, which indicates a two-phase material; the composition of
the main phase is Mn Ni Sn , and that of the secondary phase is Mn Ni Sn . The errors
47.0 41.8 11.2 68.9 28.8 2.3
of the SEM/EDX analysis are estimated with ±0.5 at%. This secondary phase, which contains a high
amount of Mn, is described as the γ-phase in the literature [41]. Liu et al. [41] reported the presence of
this phase in Ni-Fe-Co-Ga alloys and ascribed the absence of martensitic transformation in the altered
austeniticmatrixtotheγ-phase.
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a) b) c) d)
b.)
50 µm 500 µm 50 µm 500 µm
Figure 1. Optical and SEM/EDXS-micrographs of Mn Ni Sn : (a) bright field and
50 40 10
(b) back-scattered electron images; (c) bright field and (d) secondary electron images
ofMn Ni Sn .
50 41 9
In the current study, it was found that a tiny variation of the composition of approximately 1 at%
changesthemicrostructurecompletely,asshownforMn Ni Sn inthebrightfieldimageinFigure1c
50 41 9
and the secondary electron image in Figure 1d. This microstructure, with its martensitic domains and
needle-likeshapes,revealsamartensiticstructureatroomtemperature(r.t.),whichconfirmsamartensitic
transformation. For this alloy, SEM/EDXS analysis revealed a composition of Mn Ni Sn , which
48.9 41.7 9.4
isclosetothenominalcomposition.
Additionally, inductively-coupled plasma-optical emission spectroscopy (ICP-OES) analysis was
performed to verify the powder compositions. The results showed that the mean analytical
composition of the alloy Mn Ni Sn was Mn Ni Sn , which is close to the nominal
50 40 10 49.9(1) 40.2(1) 10.0(1)
composition. For the compound Mn Ni Sn , an analytical composition of Mn Ni Sn
50 41 9 49.7(1) 41.3(1) 9.0(1)
was obtained. Therefore, the analytical compositions of these alloys are in good agreement with their
nominalcompositions.
2.1.2. PhaseAnalysis
To support the results from the microstructure analysis of Mn Ni Sn , X-ray diffraction (XRD)
50 40 10
analysis was performed. Figure 2 shows the r.t. XRD patterns of the as-prepared and the post-annealed
powders. Most of the Bragg reflections of the XRD pattern of Figure 2a can be attributed to a
cubic Heusler-type crystal structure with the space group (sg) 225 (Fm3m) and a lattice parameter
a=5.997(2)Å.Thisresultindicatesanausteniticcrystalstructureatr.t.
However, in Figure 2a, a second phase could be referred to the diffraction pattern, which belongs
to the fcc γ-phase of Mn-Ni with a lattice parameter of a = 3.724(4) Å. Both determined crystal
structures support the EDXS analysis, which resulted in analytical compositions of Mn Ni Sn
47.0 41.8 11.2
andMn Ni Sn . Thisγ-phaseisknowntobepresentinNi-basedHeusleralloys,asreportedinthe
68.9 28.8 2.3
literature [41–43], appearing as a secondary phase next to the austenite. Thus, the γ phase is attributed
to a reduction of the magnetization change during the martensitic transformation and to a vanishing of
themartensitictransformation.
Metals2015,5 488

0
2
a) 2 as-prepared powder
A
b. units] A 111 A 200 111 * 200 A 311 A 222 A 400 220 A 422 311 A 440
r
a
[

I
0
sity b) A 22 post-annealed powder
n
e
t
n 2
I 1 1 0 42
A 111 A 200 * * 11* * 200 A 31 *A 222 A 40 A 331 A 420 200 *A 311 A 440
30 40 50 60 70 80 90
Scattering angle 2 [(cid:176)]
Figure 2. XRD pattern of Mn Ni Sn at r.t. for (a) as-prepared and (b) post-annealed
50 40 10
powders. “A” indicates the austenite and “γ” the fcc Mn-Ni phase. Mn-rich impurities are
labeledbyasterisks.
Afterannealingthepowderinordertoreducethestress,theBraggreflectionsshowninFigure2bare
much sharper, resulting again in a cubic Heusler-type crystal structure Fm3m with a lattice parameter
ofa=5.993(1)Åandanfccγ-phasewithalatticeparameterofa=3.711(1)Å.

1
0
a) M 1 as-prepared powder
]
s
b. unit M 100 A 110 M 110 M 002 M 200 M 112 A 211 M 211 M 202 M 103
r
a
[

I

y
t b) post-annealed powder
i
s
n Iexp
e
nt Icalc
I
Iexp - Icalc
30 40 50 60 70 80 90
Scattering angle 2 [(cid:176)]
Figure 3. XRD pattern of Mn Ni Sn at r.t. for (a) as-prepared and (b) post-annealed
50 41 9
powders. “A” indicates the austenite and “M” the martensite. Results of Le Bail fitting are
shownin(b)(R =10%).
wp
Metals2015,5 489
Concerning Mn Ni Sn (Figure 3a), the Bragg reflections are significantly different from those
50 41 9
of the Mn Ni Sn alloy, probably due to a slightly lower Sn-content. At r.t. Mn Ni Sn reveals a
50 40 10 50 41 9
martensitictetragonalcrystalstructure(P4/mmm,CuAustructuretype,sg: 123)withlatticeparameters
of a = 2.773(1) Å and c = 3.444(2) Å. The corresponding c/a-ratio is approximately 1.24. In addition
to this martensite, there is also an austenitic phase with a lattice parameter of a = 2.990(1) Å (Im3m,
tungsten structure type, sg: 229). A crystal structure with a higher order (CsCl structure type, Pm3m,
sg: 221) is indicated by a tiny Bragg reflection at 2θ = 53.72◦, which is not indexed to the tungsten
structuretype.
After recrystallization of the post-annealed powder, the diffraction pattern becomes more complex.
The diffraction pattern could be indexed to a 7M superstructure cell with lattice parameters of
a = 4.304(5) Å, b = 5.984(5) Å, and c = 42.223(7) Å and β = 93.97(2)◦, using a Le Bail fitting, which
is shown in Figure 3b. These lattice parameters are similar to those of the superstructure cell in the
Ni-Mn-Ga FMSA [44–46]. It should be noted that owing to strong reflection overlap in the Le Bail
fitting,itisnotpossibletoexcludethepresenceoftheausteniticphaseatr.t. inthissample.
2.1.3. MagneticProperties
The magnetic properties were determined by magnetometry. Figure 4a shows the
temperature-dependent magnetization analysis of the as-prepared Mn Ni Sn powder with
50 40 10
applied fields, µ H, of 0.1, 2 and 7 T in the field cooling and field heating modes, starting from
0
400 K. Concerning Mn Ni Sn at µ H = 0.1 T, three phase transformations are shown (from
50 40 10 0
low to high temperature): (i) a magnetic transformation in the martensite with a Curie temperature,
TM; (ii) a magnetostructural transformation with a martensitic start temperature, M , a martensitic
C s
finish temperature, M , and an inflection point of the martensitic transition, T ; and (iii) a magnetic
f M
transformationintheaustenitewithaCurietemperature,TA. Thus,Mn Ni Sn possessesmartensitic
C 50 40 10
transformation temperatures of M = 221 K, M = 175 K and T = 202 K and two Curie temperatures
s f M
for the martensite and austenite phases of TM = 144 K and TA = 274 K, respectively. The martensitic
C C
transformation temperatures are below r.t. Therefore, the austenite phase found at r.t. from the XRD
analysisisverified.
Additionally,themagnetizationchangeduringthemartensitictransformationis30Am2 kg−1,which
indicates a different magnetic order between the martensite and austenite. Thus, the martensite orders
ferrimagnetically,whereastheausteniteordersferromagnetically.
Themagnetizationchangeatµ H =7Tisincreasedby55%andT decreaseswithfurtherincreasing
0 M
of µ H from T = 202 K to T = 189 K, which is ascribed to a field-induced first order
0 M,0.1T M,7T
magnetostructural transformation. A reduced thermal hysteresis ∆T of 7 K, as compared to other
Hyst
Heusler alloys [11,14], was obtained, which is comparable with the results published by Han [12] and
Ma et al. [18]. However, ∆T remains constant with increasing of µ H, resulting in a possible
Hyst 0
reduction in energy consumption in magnetocaloric applications, contrary to Ni-rich Ni- and Mn-based
Heusleralloys[11,14].
Metals2015,5 490

70
]
-1 g 65 a) as-prepared powder b) post-annealed powder
2 Am k 556050 00HH == 72 TT
M [ 45 0H = 0.1 T
n 40
atio 3305
z
eti 25
n 20
g
a 15
M
10
5
0
50 100 150 200 250 300 350 400 50 100 150 200 250 300 350 400
Temperature T [K]
Figure 4. Temperature dependence of magnetization of Mn Ni Sn in applied fields
50 40 10
of µ H = 0.1, 2 and 7 T during heating and cooling modes for (a) as-prepared and
0
(b)post-annealedpowders.
Figure 4b displays the thermomagnetic curves of this alloy as a post-annealed powder using
conditions similar to those of the as-prepared powder. Comparing both at µ H = 0.1 T, the martensitic
0
transformation temperatures are very sensitive to the annealing method. Thus, M , M and T are
s f M
decreasing, whereas TA is increasing (for more details, see the Supplementary Information,Tables S1
C
and S2). Comparable results are known in Ni-rich Ni-and Mn-based Heusler alloys [19,20]. A similar
feature has been found for Ni-Co-Mn-In alloys and is attributed to the modification of the atomic sites
(order-disorder), the Mn-Mn distances and the Mn-Mn exchange coupling [47–50]. Additionally, based
on this, the Fermi surface and the Brillouin zone boundary can be affected [50–53]. Therefore, for the
post-annealed powder, the magnetization change during the martensitic transformation is decreased by
one-third and remains constant with increasing µ H up to 7 T, indicating similar magnetic order of the
0
martensiteandaustenite,i.e.,ferromagnetism.
At µ H = 0.1 T, ∆T is increased significantly by 17 K, which is comparable to Ni-rich Ni- and
0 Hyst
Mn-based Heusler alloys [11,14]. However, for Ni-rich Ni-Mn-Sn Heusler alloys reported by Schlagel,
McCallum and Lograsso [19] and Xuan et al. [20], the effect of annealing increased the magnetization
change during the martensitic transformation. This was also observed for Ni-rich Ni-Mn-Co-In Heusler
alloys, reported by Ito et al. [47]. Therefore, the impact of atomic ordering on annealing in terms of
the martensitic transformation are different between Mn-rich Ni- and Mn-based Heusler alloys because
of atomic site occupation of the Ni and Mn atoms. A recently reported study of the crystal structure
of Mn-rich Mn Ni Sn by our group [54] identified the disorder on the Wyckoff positions 4d and 4c,
1.8 1.2
byNiandMnatomsandtheireffectsonthemagneticandtransportproperties.
Concerning the temperature-dependent magnetization of Mn Ni Sn , shown in Figure 5a, a
50 41 9
small variation in composition of 1 at% increases M , M and T to higher temperatures (see the
s f M
Supplementary Information, Tables S3 and S4). Thus, at r.t., the martensite and austenite are found
to be present, verifying the crystal structures of the as-prepared powders shown in Figures 2a and 3a as
determinedbyXRDanalysisatr.t.
Comparingthemagnetizationoftheausteniteandmartensiteduringthemartensitictransformationat
µ H = 0.1 T, the differences are quite small, which means that the martensite shows paramagnetism
0
Metals2015,5 491
and the austenite orders antiferromagnetically. Thus, the main contribution during the martensitic
transformation comes from the structural distortion. At µ H = 2 T, the magnetization change reaches a
0
valueof20Am2 kg−1,showingasharpmartensitictransformation.

25
] M
-1 g a) as-prepared powder TC
2 m k 20 0H = 2 T
A 0H = 0.1 T
M [ TCM
15
zation 10 TCA TCA
eti
n
g
a 5
M
b) post-annealed powder
0
50 100 150 200 250 300 350 400 50 100 150 200 250 300 350 400
Temperature T [K]
Figure 5. Temperature dependence of magnetization analysis of Mn Ni Sn in
50 41 9
applied fields of 0.1 and 2 T during heating and cooling for (a) as-prepared and
(b)post-annealedpowders.
As shown in Figure 5b, the temperature-dependent magnetization of the post-annealed powder is
different from that of the as-prepared powder owing to a modified atomic site phenomena, which is
verifiedbythevariationofthecrystalstructuresofbothpowders(seeFigure5).
However, for the post-annealed powder shown in Figure 5b, no clear magnetization change during
the martensitic transformation was observed, indicating a similar magnetic coupling of the Mn atoms
of the austenite and martensite. Additionally, the ferromagnetic ordering of the austenite increases,
which indicates stabilization of the austenitic phase. Therefore, the martensitic transformation merges
with T of the martensite and austenite. A similar phenomenon on the magnetization curve could be
C
observed for Ni-Mn-Sn thin films [55]. The Curie temperatures of the austenite, TA, and martensite,
C
TM,increasewithincreasingappliedfieldfromµ H =0.1Tto2T(TM =153Kvs. TM =217K
C 0 C,0.1T C,2T
andTA =303Kvs. TA =331K).
C,0.1T C,2T
2.2.Co Ni Ga Nanoparticles
50 21 32
2.2.1. MicroscopicAnalysis
Figure6a,bshowstheTEMmicrographandparticle-sizedistribution,respectively,oftheas-prepared
Co Ni Ga nanoparticles. The average particle size is 120 nm, with a broad size distribution in the
50 21 32
rangeof50–200nm. AsshowninFigure6c,therearealsosomeCo Ni Ga nanorodsandnanowires
50 21 32
co-existingwiththesphericalnanoparticles.
Metals2015,5 492
a) c)
b)
Average: 120 nm
Std. Dev.: 41 nm
y
c
n
e
u
q
e
r
F
Particle size (nm)
Figure 6. (a,c) TEM micrographs and (b) particle-size distribution of the as-prepared
Co Ni Ga nanoparticlesCNG_AS.
50 21 32.5

200000 Iexp 110 100 111 200 200 211
I
calc 2
0 2
unts]150000 Iexp - Icalc 0 111 210 20 220 311 220 310
o 3
c 1
y I [100000 200 1
ntensit 50000 100 111 002 200
I
0

Scattering angle 2 [(cid:176)]
Figure 7. Experimental and the Rietveld-fitted XRD patterns of as-prepared Co Ni Ga
50 21 32
nanoparticles. Thecentral-angle2θpositionsoftheBraggreflectionsaccordingtotherefined
unit cell parameters of the phases are indicated by bars at the bottom. The difference curve
isalsoshownatthebottom. Millerindices,hkl,oftheselectedobservedreflectionsoftheβ
(B2-type ordering) and γ phases are indicated. The inset shows the scaled XRD pattern for
bettervisualizationofthefittingoftheweakreflections.
2.2.2. CrystalStructureofCo Ni Ga Nanoparticles
50 21 32
Figures 7 and 8 show the XRD patterns of CNG_AS and CNG_AQ Co Ni Ga , respectively.
50 21 32
The XRD-derivedstructural parameters(unit cellparameters, volume andweight contentof the phases)
and Rietveld fitting quality factors are compiled in Table 1. The XRD patterns exhibit an amorphous
Metals2015,5 493
halo due to the amorphous mesoporous SBA-15 silica. According to phase analysis and quantitative
Rietveldfitting,twomodificationswiththeCo NiGanominalcompositionarepresentintheas-prepared
2
nanoparticles (Figures 7). The main phase observed (94.5(1.6) wt% content) is a cubic β phase with a
B2 partial ordering (Pm3m, sg: 221, Inorganic Crystal Structure Database (ICSD) Code 169729 [35])
and a unit cell parameter of a = 2.8752(1) Å. As a minor phase with a content of 5.5(1.1) wt%, a
fully-ordered γ phase with a slightly tetragonally distorted structure (P/4mmm, sg: 123, ICSD Code
157788[34])andunitcellparametersofa=3.580(3)Åandc=3.567(5)Åisdetected(Table1). Intheβ
Co NiGastructurewithB2-typeordering,theNiandGaatomsoccupythesameposition,whereasinthe
2
tetragonal γ-phase, different atoms are placed in different positions [34,35] and the unit cell volume is
approximately two-times larger than that of the β phase. The B2-type ordering in the β phase is clearly
indicated by the presence of the weak reflection, with Miller indices hkl = 100 at 2θ = 31.12◦ and the
distinctabsenceofthesuperlatticereflectionat2θ = 26.87◦,whichischaracteristicoftheβ-phasewith
L2 atomicordering(hkl =111fortheL2 -typeβ-structure).
1 1

100000 Iexp 110 ’ 110 4
y I [counts] 6800000000 IIceaxlpc - Icalc ’ 112 0’ 004 111 200 ’ 220’ 20 211 ’ 312’ 116 ’ 224 220 ’ 008’ 400 310
ntensit 40000 100 ’ 110 ’ 20
I
20000
0
’

Scattering angle 2 [(cid:176)]
Figure 8. Experimental and Rietveld-fitted XRD patterns of annealed/quenched
Co Ni Ga nanoparticles. The central-angle 2θ positions of the Bragg reflections
50 21 32
according to refined unit cell parameters of the phases are indicated by bars at the bottom.
Thedifferencecurveisalsoshownatthebottom. Millerindiceshkloftheselectedobserved
reflections of the β (B2-type ordering) and β’ phases are indicated. The inset shows the
scaled XRD pattern for better visualization of the weak reflections. Non-indexed reflections
aremarkedbytriangles(probably112and224reflectionsoftheGa Ni phase).
7 3
As shown in Figure 8, considerable changes occur in the crystalline phase composition and content
after heat treatment. The minor tetragonal γ-phase disappears, and the content of the B2-type
ordered Co NiGa β-phase decreases to 72.2(1.4) wt%, whereas its unit cell parameter increases to
2
a = 2.8763(1) Å. The formation of a rather large amount of the highly ordered Co NiGa martensite
2
phase (β’) is observed (27.8(3.5) wt%) according to the results of the Rietveld quantitative analysis. In